Influence of Molecular Excluded Volume and Connectivity on the

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Influence of Molecular Excluded Volume and Connectivity on the Nanoscale Morphology of Conjugated Polymer / Small Molecule Blends Michael Roders, Gregory M. Pitch, Diego Garcia-Valdes, and Alexander L. Ayzner J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.7b11065 • Publication Date (Web): 27 Jan 2018 Downloaded from http://pubs.acs.org on January 28, 2018

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The Journal of Physical Chemistry

Influence of Molecular Excluded Volume and Connectivity on the Nanoscale Morphology of Conjugated Polymer / Small Molecule Blends Michael Roders, Gregory M. Pitch‡, Diego Garcia-Valdes‡, Alexander L. Ayzner* University of California Santa Cruz, Department of Chemistry and Biochemistry, 1156 High Street, Santa Cruz, CA 95064, USA. E-mail: [email protected] ‡

Authors made equal contributions.

ABSTRACT

Over the past couple of decades, organic photovoltaics (OPV) based on conjugated polymer/fullerene derivative bulk heterojunctions have been extensively studied, resulting in single-junction efficiencies of order 10%. The need to push the efficiency towards 15% has resulted in the synthesis of a large number of non-fullerene electron acceptors with ever increasing absorption coefficients in the red. Though some new acceptors have recently begun to be competitive with the fullerene, there is very little systematic understanding of which molecular geometries and spatial frontier orbital extent correlate with improved performance.

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One of the most important factors that determines the OPV efficiency is the nanoscale, phaseseparated morphology of the blend. In this manuscript, using a combination of resonant elastic X-ray scattering and elemental mapping, we investigate the influence of relatively small chemical changes to a non-planar conjugated small molecule on the nanoscale morphology of the resulting polymer/molecule blend. We find that subtle modifications of the number and placement of peripheral functional groups can have an enormous influence on the length scale of phase separation. We then quantify the extent of phase separation by using the generalized indirect Fourier transform to convert resonant scattering intensities to pair-distance distribution functions. Our results point towards the large influence of the molecular excluded volume as a major morphology determinant. This work has implications for synthetic efforts to create nonfullerene electron acceptors that can substantially outcompete fullerenes in OPV devices.

1. INTRODUCTION The discovery of the organic bulk heterojunction (BHJ) photovoltaic architecture more than two decades ago started a massive effort aimed at extracting the full potential of a molecular photovoltaic device.1-3 The main advantages of organic photovoltaics (OPV) include their inexpensive roll-to-roll fabrication, facile bandgap engineering and amenability to flexible and stretchable platforms.4 Although various combinations of polymer/polymer and smallmolecule/small-molecule5 BHJs have been studied extensively, the polymer/small-molecule device active layer has received an overwhelming share of attention. Though a myriad of novel conjugated polymers have been utilized as photoinduced electron donors, the electron acceptor small molecule was almost always a fullerene derivative.3, 6-14 Fullerenes are remarkably efficient electron acceptors;15 coupled with their quasi-isotropic intermolecular charge transfer rates,


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these molecules were deemed to be the ideal acceptor type.16 However, fullerenes suffer from relatively low molar extinction coefficients, particularly in the red and near-IR parts of the spectrum.17 Further, derivatization of the fullerene cage leads to progressive departures from charge transfer isotropy, as well as a change in electronic structure.17-19 Despite substantial polymer bandgap engineering, single-junction polymer/fullerene device conversion efficiencies appear to have plateaued in the vicinity of 10%. A clear path towards higher efficiencies is to use electron donor/acceptor combinations with increased complementary absorption, while pushing both bandgaps ever further to the red. However, the departure from the fullerene small-molecule motif generally leads to an associated loss of the quasi-isotropic charge transport.16, 20-21 To date, a great variety of non-fullerene small molecules have been used together with conjugated polymers in OPV devices. Nevertheless, the majority of these candidates have lagged behind the fullerene motif in OPV performance. In addition to charge transfer isotropy, an immensely important consideration in any novel combination of polymer and small molecule is the spatial extent of component phase separation.8, 22-24 The latter invariably takes place to varying degrees when the casting solution densifies and eventually becomes solid.23, 25-26 If the mean length scale of phase separation exceeds the relevant exciton diffusion length – a quantity that often falls in the 10 – 20 nm range – device performance suffers drastically.3, 27 We and others have shown that fullerenes often lead to phase separation on a length scale that is commensurate with the exciton diffusion length. However, there is no general understanding of how the chemical structure of a given small molecule dictates the nanoscale, nonequilibrium microstructure in the blend film. Recently, major advances in efficiencies using the ITIC (indacenodithieno[3,2-b]-thiophene)-based family


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of electron acceptors have underscored the fact that we still do not understand why this chemical structure motif results in performance improvements, even relative to the fullerene family.28 Our overarching goal is to understand how different classes of small-molecule chemical structure – and specific derivatization patterns therein – influence the nonequilibrium hierarchical microstructure in polymer/molecule blends. We began to systematically examine this question in our previous work, where we studied the coarse effect of small molecule geometry and π-electron density extent on the morphology of polymer/molecule blends.29 In this work, we extend our investigation to the amorphous spiro-bifluorene series of small-molecule derivatives in two chemically distinct, semi-crystalline polymers. The chemical structures between the derivatives differ only in the connectivity and number of identical functional groups on the molecular periphery. Neither derivative has particularly strong intermolecular interactions, leading to amorphous aggregates in all cases. This allows us to make a comparison of the hierarchical microstructure without the complication of large differences in intermolecular interactions and thus molecular packing. To characterize the microstructure on the nanometer scale, we have used a combination of elemental mapping and resonant elastic X-ray scattering. Furthermore, we have used the generalized indirect Fourier transform method to generate pair-distance distribution functions, allowing us to systematically and quantitatively interrogate the morphology of polymer blends with different small molecules. Our results point towards the importance of the molecular excluded volume as a major factor in determining the nanoscale length scale of phase separation. Further, we find that the connectivity of peripheral functional groups on the small-molecule plays a role in the nanoscale morphology. This work has implications for the current efforts to


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design small molecules that can outcompete fullerene derivatives, thereby systematically moving the non-fullerene research effort towards greater efficiencies.

2. RESULTS In this work, we explore how subtle modifications to the chemical structure of a molecular semiconductor in a BHJ thin film influences the hierarchical morphology for a given conjugated polymer and processing condition set. For this study, we focus on a chemical series of three spiro-bifluorene (Sp) derivatives (Figure 1) with varied number and configurations of peripheral phenylamino branches. The Sp series is comprised of 2,2’,7,7’-tetrakis[N,N-di(4methoxyphenyl)amino]- 9,9’-spirobifluorene (Sp-M); commonly known as Spiro-MeOTAD), 2,7-bis[N,N -bis(4-methoxy-phenyl)amino]-9,9-sprobifluorene (Sp-Z), and 2,2’-bis[N,N -bis(4methoxy-phenyl)amino]-9,9-spirobifluorene (Sp-E). The three Spiro derivatives, shown in Figure 1, share a common bridging sp3 carbon that twists the adjacent fluorene groups nearly perpendicular to one another based on molecular geometry optimizations, preventing crystallization similar to fullerene derivatives that form amorphous small-molecule domains. The choice to investigate the Sp chemical series is rooted in the fact that these molecules do not exhibit long-range order in thin films; thus, our measurements are not affected by differences in small-molecule degrees of crystallinity. Future work building on this report will systematically move beyond such considerations. We focus on equimolar blends of each Sp derivative with the archetypical semi-crystalline conjugated polymers P3HT and PTB7, and use a combination of electron microscopy and synchrotron X-ray scattering techniques to probe multiple length scales of the thin film morphology.


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Figure 1. Chemical stuctures of the Sp series (a) Sp-M , (b) Sp-Z and (c) Sp-E; chemical structures of host conjugated polymers (d) P3HT and (e) PTB7. We characterized the nanoscale structure of polymer:Sp blends in real space using scanning transmission electron microscopy energy-dispersive spectroscopy (STEM-EDS) performed at 40 kX magnification with a 200 kV accelerating voltage. This technique is suitable for mapping phase segregation in P3HT and PTB7 blends due to S atoms in the conjugated backbones of both polymers, which are absent in Sp derivatives. To probe the hierarchical morphology in our blend films in reciprocal space, we have used grazing-incidence X-ray diffraction (GIXD)30-31 and resonant elastic X-ray scattering (REXS).32-39 GIXD carries information about polymer crystallinity alone, as the small-molecules form amorphous aggregates. The combination of STEM-EDS and REXS provides a comprehensive real- and reciprocal-space nanoscale characterization on the order of 10 – 100 nm. The simultaneous use of both nanoscale probes allows us to reach robust conclusions about the dependence between Sp chemical structure and film morphology.


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2.1. Å-Scale Characterization GIXD is well-suited to studying crystallinity of thin films composed of low-Z materials. 2D-GIXD images for neat thin films of regioregular P3HT and PTB7 are shown in Figures 2a and 2b, respectively, as a function of the in-plane and out-of-plane scattering vector components ( xy and z, respectively). Thin films of P3HT are semi-crystalline, as expected,11 which is evident from three orders of (h00) Bragg reflections along the z axis and the (010) Bragg reflection peaked near the xy axis. The (100) Bragg reflection is located at = 0.38 Å-1 which corresponds to lamellar packing of d = 16.5 Å, while the (010) Bragg reflection is located at ~ 1.67 Å-1 and corresponds to a π-stacking distance of 3.76 Å. The (100) Bragg reflection in neat P3HT thin films has peak intensity on the z axis, which indicates that the polymer is primarily oriented edge-on with respect to the substrate. That is, the inter-chain π-stacking occurs largely in the substrate plane. Thin films of PTB7 are also semi-crystalline but less so than P3HT, which is evident from the broad (100) with no higher-order harmonics. The (100) Bragg reflection is located at = 0.315 Å-1 and corresponds to lamellar packing with a spacing of d = 19.9 Å, while the (010) Bragg reflection is located at = 1.58 Å-1 and corresponds to π-stacking with a spacing of d = 3.98 Å. The (100) Bragg reflection in neat PTB7 has peak intensity on the xy axis, which indicates that the polymer crystallites are oriented face-on (sidechains parallel to substrate plane) with respect to the substrate. We find that in blends of either polymer with Sp derivatives, GIXD polymer peak positions are not affected. Perhaps a more interesting question is whether the polymer crystallite orientation distribution depends on the presence and chemical structures of Sp molecules. Our


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previous work has shown that broad classes of small-molecules blended with regioregular P3HT in certain cases led to drastic modifications of the P3HT crystallite orientation distribution function (ODF, or pole figure) relative to the neat film.40-42 To investigate whether relatively subtle modifications of a given small-molecule class similarly alter the polymer ODF, the intensity of the lamellar (100) Bragg reflection was converted from Cartesian to plane polar coordinates to generate one-dimensional ODF vs. polar angle plots, shown in Figures 2c and 2d for P3HT and PTB7, respectively. The pole (0°) is normal to the film surface; the (h00) intensity


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near this region corresponds to the edge-on crystallite population. A polar angle close to 90° corresponds to the face-on fraction population. Figure 2c shows that the ODF of neat P3HT is peaked near 0°, expectedly consistent with the scattering image. Interestingly, ODFs of neat P3HT and P3HT: Sp thin films have similar peak positions and similar peak widths. P3HT:Sp-M and P3HT:Sp-E have slightly narrower ODF widths, but the broad take-away is that the P3HT ODF in blends is only minorly affected in

Figure 2. Pure polymer GIXD images films and polymer:Spiro crystallite orientation distribution functions (ODFs) for P3HT (a & c) and PTB7 (b & d) thin films. The grey bars near zero angle correspond to the forbidden region in a grazing geometry, while the bars near 90° block the VineyardYoneda wings. Both are consequences of the grazing incidence geometry.

both position and shape relative to neat films. More specifically, the ratio of face-on to edge-on crystallites as calculated from ODF integration is slightly higher for all three P3HT:Sp blends


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compared to pristine P3HT, as detailed in Table S1. The ODFs of PTB7:Sp blends are shown in Figure 2d. Compared to P3HT blends, the neat PTB7 ODF in Figure 2d indicates the presence of primarily face-on-oriented crystallites with a small population of edge-on-oriented ordered crystallites. When blended with Sp derivatives, the crystallographic texture remains primarily face-on. On the other hand, PTB7 blends have lower face-on to edge-on ratios compared to pristine PTB7, also detailed in Table S1. Thus, the polymer ODF remains qualitatively unchanged for either polymer with this small-molecule series. However, in each case the presence of small molecules moves the polymer ODF closer to isotropy, albeit from different ends of the distribution.

2.2. Current Injection Measurements Having understood the influence of each small molecule on the polymer ODF, we next ask whether there are concomitant changes to the charge transport characteristics through the polymer alone that accompany the differences in the polymer ODFs. To do this, we have used steady-state dark current injection measurements on polymer : Sp blend films in diodic architecture devices. Specifically, we have taken advantage of the fact that with the electrode configuration used herein, the current injected directly through Sp derivatives is negligible compared to the current supported by either P3HT or PTB7. We have found this to be the case with a PEDOT:PSS anode and an Al cathode. We believe this is primarily due to the mismatch between the small-molecule ionization potential with the Fermi levels of the electrodes, which could lead to formation of injection/collection barriers at the metal/organic interface. In fact, we were unable to measure meaningful currents in either pure Sp devices or in Sp blends with an insulating polymer. Thus, our changes in current as a function of small-molecule chemical


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structure are evidently overwhelmingly due to changes in the effective hole mobility of the polymer network alone. Figure 3 shows dark current density/voltage plots for as-cast polymer : Sp blends on linear and double-logarithmic scales, respectively. We note that annealing the films led to a substantial drop in current for all films, in some cases disabling the device completely. Thus, Figure 3 displays only as-cast device performance. Relative to pure P3HT films, Figure 3a shows a current decrease for P3HT:Sp-M devices by approximately a factor of 2 at 10 V, and even more so for the P3HT:Sp-Z and P3HT:Sp-E films. We observed a similar decrease in current for PTB7:Sp devices compared to pristine PTB7. Transport characteristics for both Sp-Z and Sp-E blends appear to cluster together for both P3HT and PTB7 blends. In P3HT, the Sp-M blend current slightly exceeds that of the other two derivative blends. In PTB7, this trend is reversed, as Sp-M blends lead to the smallest injected current. Upon further inspection, it is evident that the curves for the other two derivatives are no longer clustered together to the same extent as P3HT. Although the injected current as a function of bias appears comparable for the two, there is a clear slope difference, which shows that, if pushed to higher voltages, the two curves diverge. We have verified that this is in fact the case up to device breakdown, which occurs in the vicinity of 15 V. The trends illustrated in Figure 3 cannot be explained on the basis of the GIXD data alone, highlighting the importance of characterizing the blend film morphology over the entirety of its microstructural hierarchy relevant to device operation, i.e. from Å’s to 100s of nanometers.


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Figure 3. Dark Current Density-Voltage measurements of P3HT (a & b) and PTB7 (c & d) blend devices. Dashed lines in (b) and (d) correspond to power-law fits in the high-voltage region, and symbols correspond to the average measured curve. The inset shows the device thin film stack.

2.3. Nanometer-Scale Characterization of P3HT Blends We used a combination of REXS and STEM-EDS to characterize morphology over 10100s of nm – length scales relevant for exciton diffusion and long-range charge transport.14, 43 STEM-EDS elemental maps provide a semi-quantitative characterization of the nanoscale phase separation. However, EDS images are two-dimensional projections; as such, it is challenging to


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extract accurate quantitative information from the images alone. To this end, we have performed REXS experiments, which contain information about the nanoscale morphology averaged over the film thickness, with X-ray energies near the C 1s π* absorption edge (285 eV). This provides increased scattering contrast compared to conventional hard x-ray scattering techniques.39 The P3HT:Sp-M EDS map (Figure 4a) displays 200-300 nm globular regions of

Figure 4. (a)-(c) STEM-EDS Elemental Maps for (a) P3HT:Sp-M, (b) P3HT:Sp-Z, and (c) P3HT:Sp-E blends; Sulfur signal is represented by red while nitrogen signal is represented by green; (d) REXS curves modeled in GIFT including ; (e) pair distance distribution functions, , from GIFT calculations. high nitrogen (green) content embedded within regions of varying concentrations of sulfur (red), indicating the formation of large, mostly pure phase-separated Sp-M aggregates. We find that the superimposed S and N maps are clearly anti-correlated with very good agreement between the two measurements (see Supporting Information for further details). The P3HT:Sp-Z EDS map (Figure 4b) shows globular sulfur-rich polymer domains that form an interconnected


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network, as well as nitrogen-rich regions with sub-100 nm dimensions. P3HT:Sp-E blends show a qualitatively similar extent of polymer/molecule phase separation as Sp-Z blends, albeit with less defined small-molecule regions. Figure 4d shows REXS intensities as a function of the scattering vector length, . We note that in the organic solar cell literature, REXS is frequently referred to as RSoXS.38-39, 44-45 The three curves for the different Sp derivatives differ qualitatively, indicating substantial differences in the morphology. The curves have not been scaled by a monotonically increasing function of ; thus, peaks and shoulders in the scattering intensity represent physically significant correlations in real space. Unfortunately, extracting quantitative information from low- scattering experiments within a necessarily limited window is challenging. This is because structural correlations as a function of position are related to the scattering intensity in reciprocal space through the Fourier transform, which requires data sampled at all Q. This is a long-standing problem particularly in small-angle X-ray scattering. One solution to this problem appropriate to a diverse set of systems is the generalized indirect Fourier transform (GIFT) technique pioneered by Glatter.46-50 We are aware of one prior instance of the application of GIFT to polymer/molecule blends.51 However, few details about the fitting models were provided; therefore, prior results are of limited utility for understanding the range of validity and applicability of GIFT to thin blend films. We have thus undertaken the task of systematically applying GIFT to the systems described in this manuscript. To the best of our knowledge, this represents the first comprehensive attempt at using GIFT to understand the nanoscale organic semiconductor film morphology as a function of small-molecule structure. This analysis technique may be relatively unfamiliar to the organic optoelectronics community; therefore, we have summarized the main principles below prior to describing our results.


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Insofar as the scattering intensity can be written as

∝

(1)

where is the form factor of the scattering particle and is the structure factor, the GIFT algorithm allows for a model-free determination of the pair-distance distribution function (PDDF), , via

= 4

sin

(2)

Here, is an intraparticle distance that connects two points described by some electron density difference relative to its average. Such a factorization of the scattering intensity implies a model, where the sample is composed of representative (generalized) particles with an internal structure described in reciprocal space by , and whose spatial distribution in a densely packed assembly is governed by . Unlike and, by extension, , GIFT does not provide a model-free determination of . Thus, selecting a physically appropriate model is of some significance. The structure factor is related to the radial distribution function, , through the orientationally averaged Fourier transform as

= 4 − 1

sin

(3)

In GIFT, a specific multi-parameter model for is chosen and Fourier-transformed to yield the structure factor. In a particle, whose electron density fluctuations relative to the medium (or the scattering contrast) are independent of position over nanometer length scales, the PDDF simply describes the distribution of distances encountered within the particle. As such, the PDDF contains


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information about the size and shape of the particle, or the statistically representative scattering inhomogeneity. In such a case, > 0 for all . When the factorization in Equation 1 is rigorously correct, the distance at which the PDDF approaches zero may be associated with the (orientationally averaged) size of the scattering particles. When the representative scattering inhomogeneity includes regions where the scattering contrast is a function of , may change sign.50 The distance over which this occurs contains information about the average length scale that characterizes the change in intraparticle scattering contrast. In a polymer/molecule blend, this metric is related to the mean spatial phase separation extent in the interpenetrating network. The explicit assumption made in GIFT is that the scattering is isotropic, leading to a PDDF that depends only on the particle distance as opposed to vector orientation. Our results clearly show that the polymer phase contains crystalline domains with some orientational preference with respect to the substrate. This can give rise to anisotropic scattering, as others have observed.52-53 Our averaging of the scattered intensity over the polar angle, as described in the Supporting Information, thus averages away the polymer anisotropy. Though this results in some loss of information, (i) the polymer crystallite volume fraction has been previously estimated to be the minority fraction, which is further lowered in blend films, and (ii) the blend polymer crystallite ODF is not sharply peaked. Thus, we believe this kind of orientational averaging corresponds to a reasonable compromise in polymer/small molecule blends. We note that in this sample set, we did not observe anisotropic scattering. In this work, we focus on two physically plausible models for the structure factor: that of polydisperse hard spheres and mass fractal aggregates. The presence of peaks in the scattering intensity in some samples over the 0.1 to 1 nm-1 range suggests that it is reasonable to approximate the system as a dense collection of quasi-spherical particles of some average spatial


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extent. At the same time, on general grounds it may be expected that over some range of length scales, the largely disordered, highly branched and interpenetrating microstructure of a polymer/molecule blend could be associated with a fractal dimension along with its corresponding structure factor.54-56 These two structure factor models are by no means exhaustive. Substantially more intricate structure could be envisioned; however, this would be at the cost of an increase in the number of fitting parameters. Furthermore, as we show in Figures 4 and 5, we find that in practice both the polydisperse hard spheres and fractal structure factor models give rise to qualitatively – and sometimes quantitatively – similar PDDFs. However, the generality of this observation is as-yet unclear. Further details about the above structure factor models can be found in the Supporting Information. Figure 4d shows the REXS scattering intensity collected near the C 1s π* absorption edge for P3HT:Sp derivative blends corresponding to the TEM images above along with corresponding GIFT fits, which are shown as solid lines in Figure 3d. The resulting normalized PDDF plots are shown in Figure 4e. We find that for the P3HT:Sp-M blend, only the mass fractal aggregate (MFA) model for the structure factor was able to capture the details of the scattering intensity. The resulting is nonzero for all , and it is substantial even for relatively large , well in excess of the expected exciton diffusion length of P3HT. Evidently, the positivity of implies the existence of large particles with an electron density that is approximately constant on nm length scales. The distance at which the PDDF goes to zero is of order 200 nm, consistent with the STEM-EDS image in Figure 4a. Figure 4e shows that the PDDFs for P3HT blends with Sp-Z and Sp-E have a substantially different functional form compared to the Sp-M blend. The PDDF for these two derivatives undergoes a sign change – an indication that the representative “particle” or


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scattering inhomogeneity consistent with Equation 1 is composed of regions of varying scattering contrast. That is, the scattering inhomogeneity contains phase-separated regions that are enriched in P3HT or in Sp. We associate the length scale of the zero crossing with a metric related to the mean length scale of polymer/molecule phase separation. The P3HT:Sp-E blend is characterized by a with the first positive maximum at 29 nm, a negative minimum at 91 nm, and zero crossing at 59 nm. We found that, similar to P3HT:Sp-M blends, only the MFA structure factor model was able to describe the data for Sp-E. In contrast to the previous two blends, we were able to describe the P3HT:Sp-Z scattering profile satisfactorily with both the MFA and PHS models. For consistency, we have included both resultant PDDFs in Figure 4e. The MFA P3HT:Sp-Z blend PDDF exhibits the first positive maximum at 52, a negative minimum at 157 nm, and a zero crossing at 102 nm. Figure 4e shows that, although the shape of the PDDF for < 200 nm is similar between the two structure factor models, the zero crossing differs by 20 nm between the two results. Furthermore, the PHS model gives rise to an additional positive maximum at > 200 nm, a physical dimension that appears to be inconsistent with our TEM results.

2.4. Nanometer-Scale Characterization of PTB7 Blends TEM and REXS data for PTB7 are shown in Figure 5. Somewhat similar to P3HT blends, the PTB7:Sp-M sample displays 100-200 nm N-rich regions segregated from somewhat globular S-rich regions. Nevertheless, the Sp-M regions in P3HT blends appear larger and of substantially higher contrast relative to P3HT. The calculated functions using both MFA and PHS structure factor models are in agreement with the EDS map; the zero crossing is ~ 100 nm. Though both polymers share some similarities when blended with Sp-M, the shape of the PDDF


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for PTB7 differs from that of P3HT, which had > 0 for all . This again suggests less pure representative scattering inhomogeneities in PTB7 than in P3HT, consistent with the sulfur/nitrogen contrast in TEM images. We note that for this sample, PHS and MFA structure

Figure 5. (a)-(c) STEM-EDS Elemental Maps of (a) PTB7:Sp-M, (b) PTB7:Sp-Z, (c) PTB7:Sp-E; Sulfur signal is represented by red while nitrogen signal is represented by green; (d) REXS curves modeled in GIFT including ; (e) pair distance distribution functions, , from GIFT calculations. factor models gave effectively identical PDDFs. The TEM images for PTB7 blends with both Sp-E and Sp-Z show substantial intermixing as judged by the overlap of S- and N-containing regions. The qualitative trend across the derivatives is similar to P3HT; however, P3HT blends consistently gave rise to significantly more pronounced phase separation. The PDDF for Sp-Z contains a zero crossing at 48 nm, which is substantially smaller than that of Sp-M blends. For Sp-E blends, obtaining a reliable


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PDDF was complicated due to the shape of the scattering curve; the flattening of the intensity at > 0.2 nm-1 is likely related to the presence of an intensity plateau similar to the Sp-Z sample but occurring at higher outside the experimental window. As such, the PTB7:Sp-E PDDF likely has the largest uncertainty associated with it, even though the intensity fit appears to be quite satisfactory.

3. DISCUSSION Our results immediately allow us to make several qualitative statements about the microstructure of conjugated polymer/small-molecule blends. First, we have previously shown that among three diverse small-molecule classes that included planar (perylene diimide derivative), quasi-spherical (PCBM) and non-spherical/nonplanar geometries (Sp-M), some members drastically altered the conjugated polymer ODF (PCBM, perylene diimide), while SpM left it largely unchanged.29 The question remained whether it was the coarse small-molecule geometry that determined the polymer crystallite ODF, or whether fine changes to the molecule structure were sufficient to result in substantial modification of the ODF relative to the neat polymer film. The polymer ODF curves shown herein indicate that small changes to the overarching molecular motif of an amorphous small molecule, as exemplified by the spirobifluorene core, do not appear to qualitatively alter the polymer ODF. This result is supported by both P3HT and PTB7 measurements.. The two polymers in this work differed not only in their ODF but also in their relative degree of crystallinity: The PTB7 film showed only a first-order lamellar Bragg reflection, whereas P3HT showed several. On the nanometer scale, EDS maps imply that the relative contrast between the P3HT and all three small molecules is larger than the corresponding blends


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in PTB7. A possible conclusion that emerges is that with increasing polymer crystallinity, the phase-separated features become more enriched in polymer and small-molecule components, respectively. That is, polymer crystallization partially drives phase separation – an idea that has been articulated previously.22, 57 However, we cannot currently rule out the contribution that the difference in polymer molecular weights or small differences in polymer film thickness could make to this consideration. Interestingly, in P3HT:Sp-M blends with large scattering contrast and large Sp-M features, the polymer dark current density was higher than in blends with the other two derivatives. In PTB7:Sp-M blends, which showed less defined phase separation and somewhat smaller Sp-M particles relative to P3HT blends, the device current was smaller than the other derivatives. We hypothesize that with large, well-separated Sp-M particles, more P3HT chains and crystallites are on average clustered together, leading to relatively rapid charge transport between the Sp-M grains. In contrast, the still relatively large but less defined Sp-M particles in PTB7 blends present more frequent energetic impediments to transport of charges through the polymer. In other words, we speculate that when the size of the phase-separated small-molecule features becomes quite large, more well-defined phase segregation of the two components is net beneficial for the charge mobility along the polymer network. In the latter scenario, though the presence of large particles leads to effectively less volume for polymer charge transport, the conducting channels themselves are characterized by a larger effective mobility than in the case of frequent geometric impediments in a loosely segregated network. We note that this qualitative analysis of the differences in the current/voltage response of the blends relies on the implicit assumption that the different blends have at least comparable degrees of vertical phase


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segregation (if any). Unfortunately, we are currently unable to make a direct evaluation of this effect. Thus, this possibility must be born in mind. Second, the stark morphological result that was similarly robust across the two polymer hosts was that Sp-M blends consistently showed feature sizes that were well in excess of expected exciton diffusion lengths for conjugated polymers. This was not the case for Sp-E and Sp-Z blends, where the characteristic length scale of phase separation between polymer-rich and small-molecule-rich domains was more commensurate with the exciton diffusion length. Thus, although relatively small chemical changes on the Å-scale may not substantially alter the ODF, they are sufficient to induce nanoscale morphological differences. This has direct implications for the synthesis of small-molecule guests for polymer/molecule blends. Though it is evident that a much more sophisticated understanding of the link between chemical structure and morphology is required, we believe our work points to an important consideration. Within the Sp derivative series that we examined, Sp-M stood out in its large morphological features in both polymer blends. Sp derivatives here differ only in their connectivity and number of branching peripheral groups. They have nearly identical calculated angles between the sp3-linked fluorene units, so that we may reasonably expect that the molecular packing would not differ drastically among the three molecules. With these similarities in mind, combined with the common amorphous packing of all three Sp derivatives, we believe it is unlikely that the intermolecular interactions in the thin film differ substantially across this chemical series. Furthermore, at normal device blend ratios, all three Sp derivatives display comparable solubilities. These considerations suggest that the excluded molecular volume may play a role. Indeed, the solvent-excluded volume of Sp-M is 1.46 nm3 compared to 0.959 nm3 for both Sp-Z


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and Sp-E isomers. The change in this volume is not drastic, but evidently it is sufficient to change the mean length scale of phase segregation from ≤ 50 nm to > 100 nm in PTB7 blends and from ≤ 50 nm to > 200 nm in P3HT blends. To date, a great variety of non-fullerene smallmolecule electron acceptors blended into donor polymers (or donor molecules into acceptor polymers) have been explored. In both cases, the range of small molecules spans a large range of molecular excluded volumes. A reasonable argument rooted in charge transport considerations is that a small molecule characterized by an intermolecular electronic coupling that is quasi-independent of relative orientation is desired. Though promising in principle, several classes of such molecules have been unable to overtake PC60BM and PC70BM fullerene derivatives in photovoltaic efficiency.5872

To the best of our knowledge, these facts have not been reconciled rigorously in any

systematic manner. Our work suggests that modifications to the volume and/or connectivity of a given small molecule motif are important considerations when evaluating changes in the bulk heterojunction morphology and resulting device photocurrent. We do not aim to suggest that all molecules of a particular size are expected to perform poorly when blended with some conjugated polymer. Rather, within a given small-molecule motif, our data suggests that systematically increasing the excluded volume is expected to ultimately detrimentally affect the morphology. It is likely that a systematic application of GIFT across blends with a broad smallmolecule chemical series would improve our morphological understanding of non-fullerene polymer/molecule blends. Finally, we find that the molecular connectivity difference between Sp-E and Sp-Z does in fact lead to some nanoscale morphological differences. EDS maps of P3HT blends indicate that Sp-Z shows more defined phase separation, whereas Sp-E is somewhat more intermixed, in


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good agreement with the zero-crossing of the respective results. The data show that precise connectivity between certain functional groups may in some cases cause differences in polymer/small molecule intermixing. This has implications for the three-phase hypothesis, which states that the blend morphology may be described as a coexistence between pure polymer crystallite, intermixed amorphous polymer and small-molecule regions, as well as regions enriched or containing pure small-molecule domains. The thee-phase morphology has previously been associated with beneficial photovoltaic effects.14, 43 Our results suggest that changes in functional group placement may tune the molecular miscibility in the polymer phase. In PTB7 blends, the differences between the two smaller derivatives are more challenging to discern, though it is clear from the shape of the scattering curves that differences persist. Interestingly, comparing either derivative in PTB7 to its corresponding blend in P3HT shows that P3HT leads to more defined phase separation. Previously, we showed that semicrystalline (regioregular) and amorphous (regiorandom) P3HT consistently gave rise to large morphological differences. Our results indicate that, even with two semicrystalline conjugated polymers, the precise influence of the small-molecule chemical structure on the blend film morphology appears to be a function of the conjugated polymer structure.

4. CONCLUSIONS In summary, we have interrogated blends of two chemically distinct, semicrystalline conjugated polymers blended with a spiro-bifluorene small-molecule chemical series over multiple decades of length scale. In the molecular series, the connectivity and number of peripheral functional groups was varied while the Sp molar concentration and bifluorene-core motif was fixed. We found that the number of peripheral phenylamine arms, and thus the


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molecular excluded volume, was the dominant factor in determining the nanoscale microstructure of the thin film for both polymers. Further, we found more ubtle morphological differences as a function of small-molecule connectivity, the extent of which differed between the two polymers. To the best of our knowledge, this work represents the first systematic application of the indirect Fourier transform to analysis of low- scattering curves of organic semiconductor thin films. It appears that in the course of a single scattering measurement, this technique can yield quantitative information about the nanoscale morphology that is that is in good qualitative agreement with real-space STEM-EDS elemental mapping. Further systematic application of GIFT on a broad small-molecule series may allow the organic photovoltaics field to begin to rationally move beyond the fullerene derivative towards power conversion efficiencies approaching (and possibly exceeding) 15%. Supporting Information Available including experimental methods, GIXD ODF details and summary, interpretations and justifications of GIFT assumptions, structure factor model details and parameter summary, and STEM-EDS spectra and elemental maps and sulfur-void estimation.

Acknowledgements Work at the Molecular Foundry and Advanced Light Source was supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. Use of the Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, is used by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Contract No. DE-AC02-76SF00515. Molecular geometry


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optimization and solvent-excluded volume calculations were performed by Mr. Chad Townsend at UCSC.

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